For the past several years, extensive research has been devoted to the development of metal-ceramic composites, such as aluminum reinforced with carbon, boron, silicon carbide, silica, or alumina fibers, whiskers, or particles. Metal-ceramic composites with good high temperature yield strengths and creep resistance have been fabricated by the dispersion of very fine (less than 0.1 micron) oxide or carbide particles throughout the metal or alloy matrix. However, this metal ceramic composite technology has not heretofore been extended to include intermetallic matrices. Prior art techniques for the production of metal-ceramic composites may be broadly categorized as powder metallurgical approaches, molten metal techniques, and internal oxidation processes.
The powder metallurgical type production of such dispersion-strengthened composites would ideally be accomplished by mechanically mixing metal powders of approximately 5 micron diameter or less with the oxide or carbide powder (preferably 0.01 micron to 0.1 micron). High speed blending techniques or conventional procedures such as ball milling may be used to mix the powder. Standard powder metallurgy consolidation techniques are then employed to form the final composite. Conventionally, however, the ceramic component is large, i.e., greater than 1 micron, due to a lack of availability, and high cost, of very small particle size materials since their production is energy intensive, time consuming, and costly in capital equipment. Furthermore, the production, mixing and consolidation of very small particles inevitably leads to contamination at the surface of the particles. Contaminants, such as oxides, inhibit interfacial binding between the ceramic phase and the matrix, thus adversely effecting ductility of the composite. Such weakened interfacial contact can also result in reduced strength, loss of elongation, and facilitated crack propagation. In addition, the matrix may be adversely effected, as in the case of titanium which is embrittled by interstitial oxygen. Further, in many cases where the particulate materials are available in the desired size, they are extremely hazardous due to their pyrophoric nature.
Alternatively, it is known that proprietary processes exist for the direct addition of appropriately coated ceramics to molten metals. Further, molten metal infiltration of a continous ceramic skeleton has been used to produce composites. In most cases, elaborate particle coating techniques have been developed to protect the ceramic particles from the molten metal during admixture or molten metal infiltration, and to improve bonding between the metal and ceramic. Techniques such as these have resulted in the formation of silicon carbide-aluminum composites, frequently referred to as SiC/Al, or SiC aluminum. This approach is only suitable for large particulate ceramics (e.g., greater than 1 micron) and whiskers, because of the high pressures involved for infiltration. In the molten metal infiltration technique, the ceramic material, such as silicon carbide, is pressed to form a compact, and liquid metal is forced into the packed bed to fill the intersticies. Such a technique is illustrated in U.S. Pat. No. 4,444,603, of Yamatsuta et al, issued April 24, 1984. Because of the necessity for coating techniques and molten metal handling equipment capable of generating extremely high pressures, molten metal infiltration has not been a practical process for making metal-ceramic composites.
Because of the above-noted difficulties with conventional processes, the preparation of metal-ceramic composites with ceramic dispersoids in the micron size range has been extremely expensive.
Internal oxidation of a metal containing a more reactive component has also been used to produce dispersion strengthened metals, such as internally oxidized aluminum in copper. For example, when a copper alloy containing about 3 percent aluminum is placed in an oxidizing atmosphere, oxygen may diffuse through the copper matrix to react with the aluminum, precipitating alumina. This technique, although limited to relatively few systems since the two metals utilized must have a wide difference in chemical reactivity, has offered a feasible method for dispersion hardening. However, the highest possible level of dispersoids formed in the resultant dispersion strengthened metal is generally insufficient to impart significant changes in properties such as modulus, hardness, and the like. In addition, oxides are typically not wetted by the metal matrix, so that interfacial bonding is not optimum.
In recent years, numerous ceramics have been formed using a process referred to as self-propagating high-temperature synthesis (SHS), which involves an exothermic, self-sustaining reaction which propagates through a mixture of compressed powders. The SHS process involves mixing and compacting powders of the constituent elements, and igniting the green compact with a suitable heat source. On ignition, sufficient heat is released to support a self-sustaining reaction, which permits the use of sudden, low power initiation of high temperatures, rather than bulk heating over long times at lower temperatures. Exemplary of these techniques are the patents of Merzhanov et al. In U.S. Pat. No. 3,726,643, there is taught a method for producing high-melting refractory inorganic compounds by mixing at least one metal selected from groups IV, V, and VI of the Periodic System with a non-metal such as carbon, boron, silicon, sulfur, or liquid nitrogen, and locally heating the surface of the mixture to produce a local temperature adequate to initiate a combustion process. In U.S. Pat. No. 4,161,512, a process is taught for preparing titanium carbide by localized ignition of a mixture consisting of 80-88 percent titanium and 20-12 percent carbon, resulting in an exothermic reaction of the mixture under conditions of layer-by-layer combustion. These references deal with the preparation of ceramic materials, in the absence of a second non-reactive metallic phase.
U.S. Pat. No. 4,431,448 teaches preparation of a hard alloy by intermixing powders of titanium, boron, carbon, and a Group I-B binder metal, such as copper or silver, compression of the mixture, local ignition thereof to initiate the exothermic reaction of titanium with boron and carbon, and propagation of the reaction, resulting in an alloy comprising titanium diboride, titanium carbide, and the binder metal. This reference is limited to the use of Group I-B metals such as copper and silver, as binders. The process is performed with a relatively high volume fraction of ceramic and a relatively low volume fraction of metal (typically 6 volume percent and below, and almost invariably below 20 volume percent). The product is a dense, sintered material wherein the relatively ductile metal phase acts as a binder or consolidation aid which, due to applied pressure, fills voids, etc., thereby increasing density.
Another class of materials which has seen considerable interest and development is intermetallic materials, especially intermetallics of aluminum such as the aluminides of titanium, zirconium, iron, cobalt, and nickel.
The need for the advanced properties obtainable with intermetallic materials is typified by their potential application to structures capable of withstanding high temperatures, such as turbine engines. In designing and operating turbine engines today and for the foreseeable future, there are two primary problems which demand solutions from the field of materials science. The first of these is the need to operate certain portions of the engine at higher temperatures to improve operating efficiency and save fuel. The second problem is the need for lighter materials to decrease engine weight and engine operating stresses due to heavy rotating components, and to increase the operating life of disks, shafts, and bearing support structures. These latter structures require materials which are less dense than conventional nickel base superalloys, but which posses roughly the same mechanical properties and oxidation resistance as those materials in current usage.
Intermetallic compounds are particularly suited to these needs because of properties which derive from the fact that they possess ordered structures having regularly repeating (e.g., A B A B A B) atom sequencing. Modulus retention at elevated temperature in these materials is particularly high because of strong A-B bonding. In addition, a number of high temperature properties which depend on diffusive mechanisms, such as creep, are improved because of the generally high activation energy required for self-diffusion in ordered alloys.
The formation of long range order in alloy systems also frequently produces a significant positive effect on mechanical properties, including elastic constants, strength, strain-hardening rates, and resistance to cyclic creep deformation. Finally, in the case of aluminides, the resistance to surface oxidation is particularly good because these materials contain a large reservoir of aluminum that is preferentially oxidized.
However, during metallurgical processing, one problem encountered is that these materials tend to form coarse grains, which adversely effect workability, and which degrade certain mechanical properties, the most important of which is ductility. Also, in many intermetallics the strong A-B bonding results in low temperature brittleness, although the exact mechanism of the ductile-brittle transition seems to be different for the different intermetallic compounds. It is thus necessary to address the problem of minimal low temperature ductility without destroying the inherent high temperature strength and stiffness. In the prior art it has generally been considered that these latter high temperature properties may only be retained by preserving the ordered structure. However, little progress has been made in developing practical intermetallic compositions that possess sufficiently improved low temperature ductility while maintaining high temperature strength.
A conventional approach for improving the low temperature ductility of certain intermetallics involves the use of alloying additions. For example, it is known that polycrystalline Ni.sub.3 Al can be made more ductile by adding small quantities of boron. U.S. Pat. No. 4,478,791 to Huang et al describes critical composition ranges over which such boron additions are beneficial. Alloying additions have also been used to effect crystal lattice changes to produce cubic crystal structures in various intermetallics. The formation of such cubic crystal structures roduces a greater number of available slip systems within the intermetallic materials, resulting in improved room temperature ductility. The formation of cubic L12 crystal structures has been the most prefered, because the L12 structure exhibits the highest degree of symmetry of all of the possible structures for ordered alloys and therefore possesses the greatest number of avilable slip systems. The choice of suitable alloying additions is governed both by the size of the atomic nuclei and the electronic band structure of the alloying elements. An example where modification of the lattice parameter of an intermetallic material is possible is in the substitution of Al by Ti in Ni.sub.3 Al to form Ni.sub.3 (Al,Ti). Also, the substitution of Co by Ni or Fe in Co.sub.3 V has led to a series of face-centered cubic L12-type superlattices with greater ductility at ambient temperatures.
U.S. patent application Ser. No. 873,890, filed June 13, 1986, of which this application is a Continuation-In-Part, and which is hereby incorporated by reference, discloses several methods for the production of intermetallic-second phase composites. One embodiment teaches the formation of a first composite by the in-situ precipitation of second phase particles within a solvent metal matrix. This composite is then introduced into an additional metal which is reactive with the solvent metal to form an intermetallic containing matrix. The methods taught by the present invention are a modification of this process, in that an additional step is employed which consists of diluting the first composite in an additional amount of host metal to produce a second composite comprising a lower loading of the second phase particles within an intermediate metal matrix. This second composite is then introduced into another metal which reacts with the intermediate metal matrix to form an intermetallic containing matrix.